Nanostructure assisted casting of thermally stable, ultrafine grained, nanocrystalline metals

ABSTRACT

Provided herein are nanocrystalline materials comprising, e.g., a matrix including one or more metals; and nanostructures dispersed in the matrix, wherein the matrix is polycrystalline and includes grains having an average size of about 1μm or less. Also provided herein are manufacturing methods of a nanocrystalline materials.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application claims the benefit of U.S. Provisional PatentApplication No. 62/941,239, filed Nov. 27, 2019, which is incorporatedby reference herein in its entirety.

BACKGROUND

Ultrafine grained (UFG), nanocrystalline metals, exhibitingextraordinary properties, are highly demanded in the fields ofaerospace, electronics and transportation, amongst others. However,other fabrication methods for UFG, nanocrystalline metals, such asmechanical alloying, severe plastic deformation, fast cooling and thinfilm deposition, are challenging for economical mass production of bulksamples with complex geometries. Moreover, the unsatisfied thermalstability also constrains the application of UFG, nanocrystallinemetals.

It is against this background that a need arose to develop theembodiments described herein.

SUMMARY

Some embodiments include a nanocrystalline material comprising: a matrixincluding one or more metals; and nanostructures dispersed in thematrix, wherein the matrix is polycrystalline and includes grains havingan average size of about 1 μm or less. In some embodiments, the averagesize of the grains is about 600 nm or less. In some embodiments, theaverage size of the grains is about 400 nm or less. In some embodiments,the nanostructures are dispersed in the matrix at a volume fraction ofabout 5% or greater of the nanocrystalline material. In someembodiments, the volume fraction of the nanostructures in thenanocrystalline material is about 10% or greater. In some embodiments,the volume fraction of the nanostructures in the nanocrystallinematerial is about 15% or greater. In some embodiments, the matrixincludes copper, and the nanostructures include a transition metal or atransition metal carbide. In some embodiments, the matrix includes zinc,and the nanostructures include a transition metal or a transition metalcarbide. In some embodiments, the matrix includes aluminum, and thenanostructures include a transition metal carbide or a transition metalboride.

Some embodiments include a manufacturing method of a nanocrystallinematerial, comprising: heating a matrix material including one or moremetals to form a melt; loading a mixture including a salt andreinforcing nanostructures over a surface of the melt, such that thesalt is heated to form a molten salt including the nanostructuresdispersed therein; agitating the melt to incorporate the nanostructuresfrom the molten salt into the melt; delivering the melt to a molddefining a hollow space with a requisite shape; and cooling andsolidifying the melt to form a metal part including the nanocrystallinematerial and having the requisite shape.

Some embodiments include a manufacturing method of a nanocrystallinematerial, comprising: mixing a powder of a matrix material andreinforcing nanostructures to form a powder mixture; compacting thepowder mixture to form a preform; heating the preform under compressionto form a melt including the nanostructures dispersed therein; andcooling the melt including the nanostructures dispersed therein to forma master material. Some embodiments further comprise subjecting themaster material to casting to form a metal part. In some embodiments,subjecting the master material to casting includes heating the mastermaterial to form a master material melt, delivering the master materialmelt to a mold defining a hollow space with a requisite shape, andcooling and solidifying the master material melt to form the metal partincluding the nanocrystalline material and having the requisite shape.In some embodiments, cooling the master material melt is performed at arate of less than about 100 K/s.

Some embodiments include a manufacturing method of a nanocrystallinematerial, comprising: mixing a powder of a matrix material andreinforcing nanostructures to form a powder mixture; compacting thepowder mixture to form a preform; eating the preform under compressionto form a melt including the nanostructures dispersed therein;delivering the melt to a mold defining a hollow space with a requisiteshape; and cooling and solidifying the melt to form a metal partincluding the nanocrystalline material and having the requisite shape.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows microstructure of bulk UFG, nanocrystalline Cu containingdistributed WC nanoparticles. (FIG. 1A) SEM image of Cu-5 vol. % WC (bya cooling rate of about 4 K/s) acquired at about 52° showing welldispersed WC nanoparticles in the Cu matrix. Inset is the image of atypical as-cast bulk Cu-5 vol. % WC ingot with a diameter of about 50mm. (FIG. 1B) Magnified SEM image of Cu-5 vol. % WC (about 4 K/s)showing the ultrafine and nanoscale Cu grains. (FIG. 1C) SEM image ofthe cross section showing ultrafine-grained Cu matrix and WCnanoparticles present beneath surface of the sample. (FIG. 1D) A typicalFIB image of Cu-5 vol. % WC (about 4 K/s) showing the UFG,nanocrystalline microstructure. (FIG. 1E) FIB image of pure Cu castunder the same condition showing coarse Cu grains. (FIG. 1F) EBSD imageof Cu-5 vol. % WC (about 4 K/s) with grain size color code. Black phasesare WC nanoparticles, red grains are smaller than about 100 nm, yellowand orange grains are smaller than about 1.0 μm. (FIG. 1G) Summary ofthe average Cu grain sizes for different volume fractions ofnanoparticles under different cooling rates. Error bars show thestandard deviation.

FIG. 2 shows nucleation and grain growth control mechanisms bynanoparticles. (FIG. 2A) Typical DSC scanning result during the coolingof substantially pure Cu and Cu-10 vol. % WC. (FIG. 2B) A typical TEMimage of Cu-13 WC interface showing the interface between Cu and WCnanoparticle. (FIG. 2C) Fourier-filtered high resolution TEM image ofthe marked red rectangle area in (FIG. 2B) showing a characteristicinterface between WC nanoparticle and Cu matrix. Insets are the fastFourier transformation of the Cu matrix (top right) and WC nanoparticle(bottom left). (FIG. 2D) Undercooling to overcome Gibbs-Thompson pinningeffect for Cu, Al and Zn. (FIG. 2E) Schematic illustration of thenanoparticle pinning effects. (FIG. 2F) SEM image of a Cu grain refinedby WC nanoparticles. (FIG. 2G) Nanoparticles break the fundamentalconstraint in other grain refinement methods. (FIG. 2H) and (FIG. 21)Schematic illustrations of phase evolution during solidification ofsubstantially pure metal (FIG. 2H) and metal with nanoparticles (FIG.21).

FIG. 3 shows nanoparticle assisted grain refinement in other materialssystems. (FIGS. 3A-B) FIB images of Al-10 vol. % TiB2 cast by furnacecooling (about 0.7 K/s) showing the distribution of TiB2 nanoparticlesand ultrafine AL grains. (FIG. 3C) TEM image of Al-10 vol. % TiB2 (about0.7 K/s) showing one ultrafine Al grain surrounded by TiB2nanoparticles. (FIG. 3D) Al Grain size distribution of Al-10 vol. % TiB2(about 0.7 K/s). (FIGS. 3E-F) SEM image of Zn-5 vol. % WC by air cooling(about 3.7 K/s). (FIG. 3G) FIB image of Zn-5 vol. % WC (about 3.7 K/s).(FIG. 3H) Zn Grain size distribution of Zn-5 vol. % WC (about 3.7 K/s).

FIG. 4 shows thermal stability of ultrafine, nanocrystalline Cucontaining WC nanoparticles. (FIGS. 4A-D) STEM images of an area with ahigh percentage of WC nanoparticles at room temperature, about 400° C.,about 600° C. and about 850° C., respectively. (FIGS. 4E-H) STEM imageof an area with a relative low percentage of WC nanoparticles at roomtemperature, about 400° C., about 600° C. and about 850° C.,respectively. (FIG. 41) SEM image of Cu-34 vol. % WC after heattreatment (about 750° C. for about 2.0 hours). (FIG. 4J) EBSD imagecorresponds to the marked white rectangle in (FIG. 41). (FIG. 4K) Cugrain size distribution of the heat treated Cu-34 vol. % WC sample.

FIG. 5 shows fabrication of Cu containing WC nanoparticles. (FIG. 5A)Schematic illustration of the salt-assisted self-incorporation for Cu-13WC before casting bulk ingots. (FIG. 5B) Schematic illustration of thepowder melting method to cast Cu containing WC nanoparticles.

FIG. 6 shows cooling curves for furnace cooling, air cooling and waterquenching of Cu-13 WC samples.

FIG. 7 shoes the size distribution of WC nanoparticles in as-solidifiedCu-13 WC sample.

FIG. 8 shoes the structure of bulk UFG, nanocrystalline Cu containing WCnanoparticles. (A-C) FIB image of Cu-5 vol. % WC, Cu-10 vol. % WC andCu-20 vol. % WC, respectively.

FIG. 9 shows a STEM image of the WC nanoparticle rich area showing thatCu grain size is correlated with WC inter-particle spacing.

FIG. 10 shows a cooling curve during the DSC tests at a cooling rate ofabout 5° C./min.

FIG. 11 shows mechanical properties of UFG, nanocrystalline Cucontaining WC nanoparticles. (FIG. 11A) SEM image of a Cu-34 vol. % WCmicropillar machined by FIB. (FIG. 11B) FIB image of the micropillarshowing polycrystalline Cu matrix with WC nanoparticles. (FIG. 11C)Engineering stress-strain curves of as-solidified pure Cu samples(blue), with nanoparticles (black) and heat treated sample (red). (FIGS.11D-E) SEM images showing the morphology of post-deformed samples with(FIG. 11D) and without (FIG. 11E) nanoparticles. (FIG. 11F) Young'smodulus of pure Cu, Cu-19 vol. % WC and Cu-34 vol. % WC.

FIG. 12 shows an undercooling profile with respect to solid fraction.The solid curves correspond with the constitutional undercooling forAl-Ti alloys. Constrained by the phase diagram and the maximumsolubility of Ti in Al, the growth restriction factor reaches a maximumvalue of Q. in the A1-0.15Ti alloy. However, the undercooling profile ofthe nanoparticle assisted phase growth indicates an almost infinitelylarge Qn_(p).

DETAILED DESCRIPTION

Embodiments of this disclosure are directed to an improved, costeffective method to fabricate bulk, thermally stable, UFG,nanocrystalline metals through a casting process by the addition ofnanostructures (e.g., nanoparticles). UFG, nanocrystalline metals,encompassing metals with a grain size smaller than one micrometer (μm),can exhibit extraordinary mechanical, physical, and chemical properties.However, it has been considered impractical to synthesize bulk UFG,nanocrystalline metals by casting, due to its slow cooling (e.g., lessthan about 100 K/s). Moreover, other UFG, nanocrystalline metals aregenerally not thermally stable, and nanocrystalline metals (such as Al,Cu, Zn, and Mg) can exhibit extensive grain growth even at roomtemperatures. Embodiments of this disclosure can pave a path towardsscalable manufacturing of bulk, thermally stable, UFG, nanocrystallinemetals via casting. More specifically, some embodiments demonstratecasting of UFG, nanocrystalline copper (Cu), aluminum (Al) and zinc (Zn)by incorporation of nanostructures.

Aspects of some embodiments include the following:

1. Nanostructure Selection

Nanostructure selection is one consideration factor. In someembodiments, tungsten carbide (WC) nanoparticles are used for Cu and Zn,and titanium diboride (TiB2) nanoparticles and titanium carbide (TiC)nanoparticles are used for Al. The general selection standard is: (a)Nanoparticles should be thermally/chemically stable in the molten metal;(b) small molten metal-nanoparticle wetting angle (e.g., <about 90°);and (c) close lattice matching between metal matrix and nanoparticles.

2. Incorporation of Nanoparticles in Molten Metal

In some embodiments, there are at least two methods for incorporation:

(a) Melt pressing of metal powders with nanoparticles. For example: Cupowder (Sigma Aldrich, <about 10 μm) is first mixed with a designedvolume fraction of WC nanoparticles (US Research Nanomaterials, about150-200 nm) for about 1.0 hour by a mechanical shaker (SK-O330-Pro). Themixed powders are then compressed into disks of about 2.0 cm in diameterunder about 250 MPa by a hydraulic machine. Then these disks arecompression-melted at about 1250° C. by an induction heater while undera pressure of about 7-10 MPa under a graphite piston. Aftersolidification in different cooling rates (furnace cooling, air coolingand water quenching), Cu ingots with WC nanoparticles are obtained.

(b) Salt-assisted self-incorporation of nanoparticles into molten metal:nanoparticles are mixed with a salt (potassium aluminum fluoride, borax,calcium fluoride, and so forth) by mechanical mixing. A metal is meltedabove its melting temperature. Inert argon gas is purged on the surfaceof the molten metal to avoid severe oxidation. The mixedsalt-nanoparticles powder is loaded on the surface of the molten metal.Then a propeller is located below the molten metal-salt interface andoperated to stir the melt. Finally, the molten melt is allowed tosolidify to obtain ingots.

3. A Theoretical Model of Nanostructure Assisted Continuous Nucleationand Controlled Grain Growth

A theoretical model is successfully established to explain the castingof UFG, nanocrystalline metals. Several mechanisms can contribute: (a)Large amount of nanostructures can serve as potential nucleation sitesduring solidification; (b) continuous nucleation during solidification;(c) nanostructures impede the grain growth by Gibbs-Thompson effects;and (d) nanostructures tune the thermal properties (e.g., thermalconductivity) of the melt.

4. Extraordinary Termal Sability of UFG, Nanocrystalline Metals Assistedby Nanostructures

As-cast UFG, nanocrystalline Cu with WC nanoparticles shows outstandingthermal stability at high temperatures. The UFG structure is stable atabout 750° C. and has little grain growth at about 850° C.

Advantages of embodiments of this disclosure include:

Casting process can be used to achieve casting of bulk UFG,nanocrystalline metals.

Scalable manufacturing process.

Excellent thermal stability up to about 750° C. or greater.

Excellent mechanical properties.

EXAMPLE EMBODIMENTS

In some embodiments, a nanocrystalline material includes a matrixincluding one or more metals, along with reinforcing nanostructuresdispersed in the matrix. Examples of suitable matrix materials includealuminum (Al), magnesium (Mg), iron (Fe), silver (Ag), copper (Cu),manganese (Mn), nickel (Ni), titanium (Ti), chromium (Cr), cobalt (Co),zinc (Zn), alloys, mixtures, or other combinations of two or more of theforegoing metals, and alloys, mixtures, or other combinations of one ormore of the foregoing metals with other elements. In some embodiments,the matrix is polycrystalline and includes grains having an average size(or an average dimension) of about 2 μm or less, about 1.5 μm or less,about 1.2 μm or less, about 1 μm or less, about 800 nm or less, about600 nm or less, about 400 nm or less, or about 300 nm or less, and downto about 200 nm or less, or down to about 150 nm or less.

In some embodiments, the nanostructures can have at least one dimensionin a range of about 1 nm to about 1000 nm, such as about 1 nm to about500 nm, about 1 nm to about 400 nm, about 1 nm to about 300 nm, about 1nm to about 200 nm, or about 1 nm to about 100 nm. In some embodiments,the nanostructures can have at least one average or median dimension ina range of about 1 nm to about 500 nm, about 1 nm to about 400 nm, about1 nm to about 300 nm, about 1 nm to about 200 nm, or about 1 nm to about100 nm. In some embodiments, the nanostructures can includenanoparticles having an aspect ratio of about 5 or less, or about 3 orless, or about 2 or less and having generally spherical or spheroidalshapes, although other shapes and configurations of nanostructures arecontemplated, such as nanofibers and nanoplatelets. In the case ofnanoparticles of some embodiments, the nanoparticles can have at leastone dimension (e.g., an effective diameter which is twice an effectiveradius) or at least one average or median dimension (e.g., an averageeffective diameter which is twice an average effective radius) in arange of about 1 nm to about 1000 nm, such as about 1 nm to about 500nm, about 1 nm to about 400 nm, about 1 nm to about 300 nm, about 1 nmto about 200 nm, or about 1 nm to about 100 nm.

In some embodiments, the nanostructures can include one or moreceramics, although other nanostructure materials are contemplated, suchas metals. Examples of suitable nanostructure materials include metaloxides (e.g., alkaline earth metal oxides, post-transition metal oxides,and transition metal oxides, such as aluminum oxide (Al₂O₃), magnesiumoxide (MgO), titanium oxide (TiO₂), and zirconium oxide (ZrO₂)),non-metal oxides (e.g., metalloid oxides such as silicon oxide (SiO₂)),metal carbides (e.g., transition metal carbides, such as titaniumcarbide (TiC), niobium carbide (NbC), chromium carbide (Cr₃C₂), nickelcarbide (NiC), hafnium carbide (HfC), vanadium carbide (VC), tungstencarbide (WC), and zirconium carbide (ZrC)), non-metal carbides (e.g.,metalloid carbides such as silicon carbide (SiC)), metal silicides(e.g., transition metal silicides, such as titanium silicide (TiSi)),metal borides (e.g., transition metal borides, such as titanium boride(TiB₂), zirconium boride (ZrB₂), hafnium boride (HfB₂), vanadium boride(VB₂), and tungsten boride (W₂B₅)), metal nitrides (e.g., transitionmetal nitrides), non-metal nitrides (e.g., metalloid nitrides such assilicon nitride), alloys, mixtures, or other combinations of two or moreof the foregoing. Particular examples of suitable nanostructurematerials include transition metal borides (e.g., TiB₂) and transitionmetal carbides (e.g., TiC and WC), amongst other transitionmetal-containing ceramics.

In some embodiments, the nanocrystalline material can include thenanostructures at a high volume percentage of, for example, greater thanabout 3%, such as about 5% or greater, about 6% or greater, about 7% orgreater, about 8% or greater, about 9% or greater, about 10% or greater,about 15% or greater, about 20% or greater, about 25% or greater, orabout 30% or greater, and up to about 40% or greater.

In some embodiments, the matrix includes Cu, and the nanostructuresinclude a transition metal (e.g., W) or a transition metal carbide(e.g., WC).

In some embodiments, the matrix includes Zn, and the nanostructuresinclude a transition metal (e.g., W) or a transition metal carbide(e.g., WC).

In some embodiments, the matrix includes Al, and the nanostructuresinclude a transition metal carbide (e.g., TiC) or a transition metalboride (e.g., TiB₂).

In some embodiments, a manufacturing method of a nanocrystallinematerial includes: (1) heating a matrix material including one or moremetals to form a melt; (2) loading a mixture including a salt andreinforcing nanostructures over a surface of the melt, such that thesalt is heated to form a molten salt including the nanostructuresdispersed therein; (3) agitating the melt to incorporate thenanostructures from the molten salt into the melt; and (4) cooling themelt including the nanostructures dispersed therein to form a mastermaterial.

In some embodiments of the method, the method further includessubjecting the master material to casting to form a metal part. In someembodiments, subjecting the master material to casting includes heatingthe master material to form a melt, delivering the melt to a molddefining a hollow space with a requisite shape, and cooling andsolidifying the melt to form the metal part including thenanocrystalline material and having the requisite shape. In someembodiments, cooling the melt is performed at a rate of less than about100 K/s, such as about 90 K/s or less, about 70 K/s or less, about 50K/s or less, about 30 K/s or less, or about 10 K/s or less.

In some embodiments of the method, features of the matrix material, thenanostructures, and the nanocrystalline material are as described forthe foregoing embodiments of the nanocrystalline material.

In some embodiments, a manufacturing method of a nanocrystallinematerial includes: (1) mixing a powder of a matrix material andreinforcing nanostructures to form a powder mixture; (2) compacting thepowder mixture to form a preform, such as in a hydraulic press and undera pressure of about 100 MPa to about 300 MPa, or about 250 MPa; (3)heating the preform under compression, such as under a pressure of about7 MPa to about 10 MPa and to a temperature up to or below (e.g., up toabout 100° C. below, or up to about 200° C. below) a melting temperatureof the matrix material, to form a melt including the nanostructuresdispersed therein; and (4) cooling the melt including the nanostructuresdispersed therein to form a master material.

In some embodiments of the method, the method further includessubjecting the master material to casting to form a metal part. In someembodiments, subjecting the master material to casting includes heatingthe master material to form a melt, delivering the melt to a molddefining a hollow space with a requisite shape, and cooling andsolidifying the melt to form the metal part including thenanocrystalline material and having the requisite shape. In someembodiments, cooling the melt is performed at a rate of less than about100 K/s, such as about 90 K/s or less, about 70 K/s or less, about 50K/s or less, about 30 K/s or less, or about 10 K/s or less.

In some embodiments of the method, features of the matrix material, thenanostructures, and the nanocrystalline material are as described forthe foregoing embodiments of the nanocrystalline material.

EXAMPLES

The following example describes specific aspects of some embodiments ofthis disclosure to illustrate and provide a description for those ofordinary skill in the art. The example should not be construed aslimiting this disclosure, as the example merely provides specificmethodology useful in understanding and practicing some embodiments ofthis disclosure.

Bulk Ultrafine Grained, Nanocrystalline Metals via Slow CoolingOverview:

Cooling, nucleation, and phase growth are ubiquitous processes innature. Effective control of nucleation and phase growth is ofsignificance to yield refined microstructures with enhanced performancefor materials. Ultrafine grained (UFG), nanocrystalline metals canexhibit extraordinary properties. However, other microstructurerefinement methods, such as fast cooling and inoculation, have reachedcertain fundamental constraints. It has been considered impractical tofabricate bulk UFG, nanocrystalline metals via slow cooling. Here thisexample reports that nanostructures can refine metal grains to ultrafinenanoscale by instilling a continuous nucleation and growth controlmechanism during slow cooling. The bulk UFG, nanocrystalline metal withnanostructures also reveals an unprecedented thermal stability. Thismethod overcomes the grain refinement constraints and can be extended toany other processes that involve cooling, nucleation and phase growthfor widespread applications.

Introduction:

Cooling, nucleation, and phase growth are ubiquitous processes ofsignificance in various aspects, such as cloud formation, icenucleation, and volcanic rock evolution. It is established thateffective control of nucleation and phase growth will yield refinedmicrostructures with enhanced performance for materials, and hence vitalto numerous broad fields, including materials science, climate andatmospheric sciences, biomedicine and chemistry. Widely usedtechnologies that involve cooling, such as casting, are of significancefor the mass production of complex materials and components. Ultrafinegrained (UFG), nanocrystalline metals can exhibit extraordinaryproperties. However, it has been considered impractical to fabricatebulk ultrafine grained, nanocrystalline metals by casting, partly due toits slow cooling (e.g., less than about 100 K/s). Other microstructurerefinement methods, such as fast cooling (thousands to millions K/s) andinoculation, have reached certain fundamental or technical constraints.Fast cooling substantially restricts the size and complexity ofas-solidified materials. The minimum grain size generally achievable byinoculation in casting falls in the range of tens of micrometers.

Grain refinement in metals during solidification is of great interestdue to the enhanced mechanical properties, more homogeneousmicrostructure and improved processability of refined microstructures.Various approaches such as inoculation, growth restriction by addingalloy elements and fast cooling (up to about 10⁷ K/s) are investigatedthrough both theoretical and experimental pathways in an effort toattain optimal grain refinement effects. However, these approaches havefailed to demonstrate whether it is possible to cast UFG,nanocrystalline metals via a casting process, which would represent arevolutionary approach given the pervasiveness of casting inmanufacturing. As a consequence of the inability to implementsolidification processes to fabricate UFG, nanocrystalline metals,various methods emerge for the fabrication of UFG, nanocrystallinemetals, including: mechanical alloying, severe plastic deformation andthin film deposition and although some degree of success has beenachieved using these processes they remain constrained to the solidstate and present challenging issues for economical mass production ofbulk samples with complex geometries. Solidification behavior of metalscan be controlled by the addition of nanostructures.Nanostructure-controlled solidification can provide a pathway forcasting metals with refined microstructures.

Here this example reports that nanostructures can refine metal grainsdown to ultrafine or even nanoscale by instilling a continuousnucleation and growth control mechanism during slow solidification. Whencasting substantially pure copper (Cu) with tungsten carbide (WC)nanoparticles, the grain sizes of Cu are refined substantially down toultrafine and even nanoscale. The as-solidified bulk ultrafine,nanocrystalline Cu reveals an unprecedented thermal stability up toabout 1023 K (0.75 melting point of Cu) and high mechanical properties.Furthermore, this revealed grain control mechanism is successfullyapplied in other materials systems such as aluminum-titanium boride(Al-TiB₂) and zinc-tungsten carbide (Zn-WC) for ultrafine grains viaslow cooling. This method paves a pathway for the mass production ofbulk stable UFG, nanocrystalline materials that can be readily extendedto any other processes that involve cooling, nucleation and phase growthfor widespread applications.

Results and Discussion: Nanoparticles Incorporation and Dispersion inBulk Samples

Preparation is made of bulk Cu ingots with WC nanoparticles by twodifferent methods. The first method is a salt-assistedself-incorporation of nanoparticles into molten metal (see Methods). Asshown in FIG. 5A, molten salt (Borax+about 5% CaF₂) could dissolve theoxide layer at the top of molten metal and provide a clean interfacebetween Cu melt and nanoparticles. In addition, the wetting anglebetween Cu and WC at about 1250° C. is below about 10°, which indicatesa good wettability between Cu and WC so that WC nanoparticles prefer totransport from the molten salt to the Cu melt to reduce the systemenergy. Combined with mechanical mixing, WC nanoparticles can be readilyincorporated into molten Cu. This salt-assisted incorporation methodopens up a scalable manufacturing method to fabricate metals containingdifferent volume percentages of nanoparticles. In this example, bulk Cuingots with about 5%, about 10% and about 20% volume fraction of WCnanoparticles were cast. Another method is a powder melting process (seeMethods) especially suitable for high volume percent of nanoparticles asshown in FIG. 5B. A cold compacted Cu-13 WC powder preform was melted byan induction heater under a pressure of about 7-10 MPa. Cu ingots withabout 19% and about 34% volume fraction of WC nanoparticles werefabricated by this powder melting method. It should be noted that thesecond method may not be as scalable as the first method, but is moresuitable for a high volume loading of nanoparticles if desired.

To evaluate the effects of the cooling rate under a same initialcondition, all these Cu-13 WC samples were then melted again and castunder different cooling rates, namely furnace cooling (about 2-4 K/s),air cooling under the protection of argon gas (about 7-12 K/s) and waterquenching (about 70-100 K/s). The typical cooling curves are shown inFIG. 6, and the thermal arrest caused by the solidification of Cu isclearly identified in the furnace and air cooling curve, while not sonoticeable in the water quenching curve due to the rapid heatdissipation.

Grain structure characterization

Characterization is made of the distribution and dispersion of WCnanoparticles and grain structures in cast bulk Cu samples via scanningelectron microscope (SEM), transmission electron microscope (TEM),focused ion beam (FIB) imaging, scanning transmission electronmicroscope (STEM) and electron backscatter diffraction (EBSD). SEMsamples were sectioned along a cross-section of the sample. To clearlyreveal the nanoparticles, the mechanically ground and polished sampleswere further polished by low-angle ion milling at about 4° for about 1.5hours.

The inset in FIG. 1A shows one typical bulk Cu-13 WC ingot cast afterthe salt-assisted incorporation method. FIG. 1A shows the typical SEMmicrostructure of Cu-5 vol. % WC by furnace cooling acquired at about52° . WC nanoparticles are uniformly dispersed in the Cu matrix. Atheoretical analysis, detailed in the subsection below entitled‘Nanoparticle dispersion and self-stabilization mechanism’, attributesthat the excellent wettability between Cu and WC can stabilize thedispersed nanoparticles in the metal melt. The average size of WCnanoparticles was measured to be about 200 nm in diameter as shown inFIG. 7. The UFG microstructure of Cu-5 vol. % WC (about 4 K/s) is shownin the SEM image of FIG. 1B (Cu grains marked by white dash lines).Grains in this SEM image are smaller than about 1000 nm. It should benoted that some areas without nanoparticles from a top view under SEMmay have nanoparticles beneath the surface as shown in the SEM image ofFIG. 1C, showing the cross-section cut by FIB. Cu grains are pinned bynanoparticles beneath the surface. Pt coating is used to protect the Cusurface during the FIB cutting. The channeling contrast of differentgrains induced by ion beam makes FIB a powerful tool to characterizegrain structures. FIG. 1D and FIG. 1E are the typical ion beammicrographs of Cu-5 vol. % WC and substantially pure Cu cast under thesame casting condition (about 2-4 K/s), respectively. The dark phase inFIG. 1D corresponds to WC nanoparticles whereas the white or grayishphases indicate the Cu grains. More FIB images of as-cast Cu-13 WCsamples with about 5, about 10 and about 20 vol. % of WC nanoparticlesare shown in FIGS. 8A-C. With the addition of WC nanoparticles into Cu,the grain sizes are readily refined to ultrafine and even nanoscale.From the FIB micrographs, it can be observed that a higher percentage ofnanoparticles yield more dense nanoparticles distribution and thus morerefined grain microstructures. For comparison, the as-solidified pure Cusample, however, has an average grain size of 270±132 μm under the samecooling rate.

EBSD analysis was used to further investigate the grain sizes. FIG. 1Fshows the typical EBSD micrograph of Cu-5 vol. % WC cast by a coolingrate of about 4 K/s. The black phases are corresponding to the WCnanoparticles. Given that the surfaces of nanocomposite samples were notperfectly flat (nanoparticles stick out from the matrix) after ionmilling, some regions (marked as greyish region) were not identifiableduring the EBSD scanning. The colors of different grains correspond todifferent sizes, as shown in the legend. Red grains are smaller thanabout 100 nm, yellow and orange grains are smaller than about 1.0 μm,while green and deep blue grains are smaller than about 2.0 μm. Themajority of the Cu grains are smaller than about 1.0 μm and significantnumber of nanosized Cu grains (marked with red color in the EBSDmicrograph) can be observed in the nanoparticle-rich areas. The averagegrain sizes of different areas under different cooling rates aresummarized in FIG. 1G. The average grain size of Cu-13 WC samples withdifferent fraction of WC nanoparticles ranges from about 236-434 nm,about 227-384 nm and about 193-347 nm for furnace cooling, air coolingand water quenching, respectively. The error bars indicate the standarddeviation of the grain size measurements. It is shown that a higherpercentage of nanoparticles yields more refined grains. FIG. 1G alsoindicates that the grain size slightly decreases with an increasedcooling rate although the differences from the cooling rates are not assignificant. The EBSD scanning also supports that the addition of WCnanoparticles can refine Cu grains to ultrafine/nanoscale by regularcasting.

If it is assumed that nanoparticles are spherical and homogeneouslydistributed in the matrix, the theoretical inter-particle spacingbetween WC nanoparticles can be calculated by:

$d = {r\left( \frac{4}{3f_{v}} \right)}^{\frac{1}{3}}$

where d is the theoretical inter-particle spacing, r is the radius ofthe particles (e.g., about 100 nm in this case), and f_(v) is the volumefraction of particles. The green dotted line in FIG. 1G corresponds tothe theoretical inter-particle spacing of the WC particles at differentvolume fractions. It is noted that the experimental results for Cu grainsizes are close to the theoretical inter-particle spacing betweennanoparticles. FIG. 9 shows the scanning transmission electronmicroscope (STEM) image of the nanoparticle rich area, which indicatesthat Cu grain size is correlated to the WC inter-particle spacing. Theequation also indicates that a smaller nanoparticle would allow a lowervolume fraction to yield smaller grains as long as the nanoparticles arenot engulfed. Lower cooling rates actually favor the non-engulfment(pushing) of nanoparticles during solidification.Nucleation and grain growth control

To further provide fundamental insight into the underlying nucleationand grain growth phenomena, differential scanning calorimetry (DSC)analyses were conducted and the typical cooling curves at a cooling rateof about 5° C./min of substantially pure Cu and Cu with WC nanoparticles(samples have the same mass, about 50 mg) are shown in FIG. 2A (see moredetails in Methods and FIG. 10). The bump on the cooling curve ofsubstantially pure Cu sample corresponds to the exothermal peak from thesolidification of Cu started at about 1033° C., which indicates an about51° C. undercooling to activate major nucleation of Cu grains. However,the starting point of the exothermal peak of Cu-10 vol. % WC is about1078° C., which means about 6 ° C. undercooling results by the additionof WC nanoparticles. The undercooling difference indicates that WCnanoparticle is a relatively potent nucleation particle for Cu. Foreffective nucleation, the crystallographic lattice discrepancy betweenparticles and matrix is of significance for the nucleation; thus theinterface between matrix and nanoparticles were evaluated at atomicscale by high resolution transmission electron microscope (HRTEM). FIG.2B shows the typical interface between WC nanoparticles and Cu matrix atatomic scale. FIG. 2C is the Fourier-filtered atomic resolution TEMimage at the marked area of FIG. 2B. The atomic structure indicates aclean and well-matched Cu-13 WC interface with no intermediate phasespresent. The top-right and bottom-left insets are the fast Fouriertransformation of the Cu matrix and WC nanoparticles, respectively. Itis identified that the (1011) planes of WC nanoparticles are parallelwith (200) planes of Cu matrix. The plane distance of (200) Cu and(1011) WC are about 0.1806 nm and about 0.1881 nm, respectively. Thusthe misfit is calculated to be about 4.1%, which implies a coherentlattice matching at this specific interface. The presence of a clean andcoherent interface indicates a strong interfacial bond. The coherentinterface indicates that WC nanoparticles could serve as potentnucleation sites for Cu grains.

Moreover, the number of potential nucleation sites (WC nanoparticles) inCu-13 WC sample is significantly higher than that in an inoculationmethod. Nanoparticles can serve as effective nucleation sites due totheir high number density. For example, the population density ofparticles for 5 parts per thousand (ppt) Al-5Ti-1B refiner is about5×10¹⁴ m⁻³. In comparison, the population densities of WC nanoparticlesare calculated to be about 1.2×10¹⁹m⁻³, about 2.4×10¹⁹m⁻³ and about4.8×10¹⁹m⁻³ (assuming that WC nanoparticles are spherical) for Cu-5 vol.% WC, Cu-10 vol. % WC and Cu-20 vol. % WC samples, respectively. Thepopulation densities of nanoparticles are five orders of magnitudehigher than that with the Al—Ti—B grain refiner.

Following the initial nucleation of grains from the melt, the nucleatedgrains grow rapidly and release latent heat, which usually impedes thenucleation of additional grains nearby. In contrast, with nanoparticlesin the molten metal, nanoparticles can rapidly assemble/adhere to thesolid-liquid interface to effectively restrict the grain growth,preventing other potential nucleation sites from being suppressed by thelatent heat release. Therefore, later nucleation events can occurcontinuously during the solidification, which is significant for thisgrain refinement and growth control mechanism. This growth restrictionand continuous nucleation can be validated by comparing the width of theexothermal peaks of substantially pure Cu and Cu with nanoparticles inDSC analyses. FIG. 2A indicates that the Cu-10 vol. % WC sample tookabout 83% longer time to complete the solidification. The heat flowcurves during DSC tests are shown in FIG. 10. The exothermal peakscorrespond to the latent heat release from the solidification. The widthof the peak in substantially pure Cu is about 6.9° C., while it is about12.6° C. in the Cu-10 vol. % WC sample, which means exothermal peak insubstantially pure Cu is sharp and intensive, and it is gradual and widein Cu-13 WC samples. These results indicate that the nanoparticles areable to slow down solidification and allow a continuous nucleation andan effective grain growth control during solidification.

After nucleation, grain growth control during solidification is ofsignificance to achieve refined grain structures. It is proposed herethat nanoparticles will impede grain growth during solidification byforcing the solidification front to grow with a non-linear geometry(e.g., curved solidification front). When Cu grains form curvatures witha small radius during solidification, the free energy (Gibbs-Thomsoneffect) increases. The molar free energy increase can be calculated by:

${\Delta G_{\gamma}} = \frac{2\gamma V_{m}}{r}$

where ΔG_(γ) is the molar free energy increase, γ is the interfacialenergy of the interface, V_(m) is the molar volume of the phase, and ris the radius of curvature of the interface. If this increment of freeenergy is large enough (thus larger undercooling is involved), thegrowth of Cu grains may be inhibited. The undercooling to overcome acurved interface generated by nanoparticles pinning can be described bythe Gibbs-Thompson effects and can be calculated by:

${\Delta T} = \frac{2\gamma T_{m}}{rH_{f}}$

where y is the interface energy of solid-liquid interfaces (Cu: about0.185 J/m², Al: about 0.116 J/m² and Zn: about 0.09 J/m²), T_(m) ismelting of the metal, r is the radius of curvature of the solidificationfront and H_(f) is the enthalpy of fusion. The undercooling to overcomethe grain front curvatures when pinned by nanoparticles for Cu, Al, andZn are shown in FIG. 2D.

As shown in the schematic illustration (FIG. 2E) of the nanoparticlepinning effects, there are mainly two types of curvature generated bynanoparticles: (1) Marked as r₁ and r₃ in FIG. 2E, when a solidificationfront meets with dispersed nanoparticles, the Cu phase can grow throughthe micro/nano-channels between nanoparticles and form a curvature witha small radius that results in an increase of free energy (Gibbs-Thomsoneffect). (2) Marked as r₂ in FIG. 2E, nanoparticles have curved surfacesand when Cu solidification front meet nanoparticles, the curvature atthe interface between Cu and WC nanoparticles will involve extraundercooling to remain solid. The range of the undercooling specified bythe curvatures generated by nanoparticle pinning effects areapproximately in the range of tens degrees of Kelvin as marked by r₁, r₂and r₃ in FIG. 2D. The curvatures generated by nanoparticles areobserved in SEM image as shown in FIG. 2F. Cu grain is surrounded byseveral WC nanoparticles and shows both types of curvatures as mentionedabove.

On the basis of theories for nucleation and grain growth in alloys,grain growth velocity can be reduced by the solute atoms at thesolid/liquid interface during alloy solidification. An empiricalrelationship (as shown in FIG. 2G) can be revealed between average grainsize (d) and the restriction factor Q for alloys:

$d = {a + \frac{b}{Q}}$

where a is a constant related to the number of particles that actuallynucleate grains at infinite values of Q. In the ideal case, when Q isinfinity large, the grain size could be refined to a which is dictatedby inter-particle spacing. b is a constant related to the potency of thenucleation particles. Q is the restriction factor, which is inverselyproportional to the constitutional undercooling, traditionally specifiedas:

Q=mC ₀(k−1)

where m is the liquidus slope in a linear phase diagram, C₀ is thesolute content in the alloy and k is the equilibrium solute partitioncoefficient. From the equation of Q, it can be seen that the restrictionfactor is determined by the phase diagram and the specific chemicalcomposition of the specific alloy. Thus the chemical restriction fromsolute atoms is constrained. The maximum value of Q (Q_(max)) in somepractical alloys can be about 50 K. As shown in FIG. 2G, with aconstrained Qmax, the smallest grain size achievable by inoculation andsolute atoms is d_(limit), which is approximately tens of micrometers.

In contrast, here it is proposed that nanoparticle pinning can induce arestriction factor, Q_(np), which can break the fundamental constraintset by Q that depends on constitutional undercooling. The nanoparticleassisted grain growth restriction can extend the chemical restrictionfactor Q to a physical restriction factor Qn_(np), which could bereadily increased to a significant large number, if not infinity (seebelow section entitled ‘Growth restriction factor by nanoparticles’ andFIG. 12). Moreover, populous nanoparticles (five orders of magnitudemore) could serve as potent nucleation sites in the continuousnucleation; thus a (related to the number of effective nucleation sites)is significantly reduced. The average grain size can be reduced to dminwhen 1/Q_(np) approaches to zero. It is proposed that d_(min) achievableis determined by the theoretical inter-particle spacing. The trend lineof theoretical inter-particle spacing in FIG. 1G is consistent with theexperimental data of average grain sizes in Cu-13 WC samples. Thisindicates that nanoparticle assisted continuous nucleation and graingrowth control can refine grains down to the inter-particle spacingunder slow cooling. Therefore, nanoparticles effectively break thefundamental constraints set by other solidification processes andreadily refine grains down to ultrafine or even nanoscale by regularcasting process. Based on the experimental results and theoreticalanalysis, the mechanisms of nanoparticles assisted grain refinement andcontrol, in comparison to other grain refinement, during regular castingare schematically illustrated in FIG. 2H and 2I.

Other less dominant mechanisms may also play minor roles in phase growthcontrol, such as blocking of the diffusion of atoms to the surface ofthe growing phase and modification of the local temperature field bynanoparticles. Nanoparticles can remain at the solidification front andblock the transportation of the atoms, thus slowing down the graingrowth. When the density of nanoparticles is low, this effect will notplay a major role. Moreover, when WC nanoparticles are close to thesolidification front, the nanoparticles could affect the local thermalfields. The lower thermal conductivity of ceramic nanoparticles couldslow down the transportation of latent heat from the solidificationfront and protect potential nucleation sites.

Other Materials Systems

The nanoparticle assisted grain refinement has also been validated forother materials systems. Al—TiB₂ and Zn—WC were also evaluated in orderto determine whether it is possible to achieve UFG, nanocrystallinemicrostructures via slow cooling for different metals. The salt-assistedself-incorporation method (see Methods) was used to fabricate Alcontaining about 10 vol. % TiB₂ nanoparticles (with an average size ofabout 100 nm) by furnace cooling (about 0.7 K/s as measured). FIG. 3Ashows the FIB image of Al-10 vol. % TiB₂, where the dark phases are TiB₂nanoparticles while the grey matrix is Al. TiB₂ nanoparticles arereasonably well dispersed in the Al matrix. The higher magnified imageis shown in FIG. 3B, and one can clearly identify Al grains (marked bywhite dash lines and pink arrows) smaller than about 1.0 μm surroundedby TiB₂ nanoparticles (marked by yellow arrows). TEM image of FIG. 3Cdemonstrates one ultrafine Al grain (marked by white dash lines)surrounded by several TiB₂ nanoparticles. The grain size distribution ofAl is shown in FIG. 3D with an average grain size of 460±220 nm. Zn-5vol. % WC samples were also fabricated by casting with a cooling rate ofabout 3.7 K/s (see Materials and Methods). As shown in FIG. 3E and 3F(SEM images of Zn-5 vol. % WC), WC nanoparticles are distributed anddispersed in Zn matrix. Zn grains (marked by white dash lines and pinkarrows) close to about 1.0 μm are clearly identified. The FIB image ofFIG. 3G shows a better contrast from the Zn grains. The sizedistribution of Zn grains is shown in FIG. 3H with an average grain sizeof 991±746 nm. It is thus validated that Zn grains were refined to about1.0 μm by the addition of 5 vol. % WC nanoparticles, although not aseffective as in the Cu and Al cases. The reason is mainly due to thesmaller undercooling to overcome the Gibbs-Thompson pinning effects forZn than for Al and Cu, as shown in FIG. 2C. This nanoparticle assistedgrain refinement approach provides an additional general pathway toproduce UFG, nanocrystalline metals by casting.

Thermal Stability

The thermal stability of UFG, nanocrystalline metals is another grandchallenge that constrains their widespread use in many applications.Pure nanocrystalline metals (e.g., Al, Cu, Sn, Pb, Zn, and Mg) canexhibit extensive grain growth even at room temperature (about 300 K).Nanocrystalline metals with higher melting points (e.g., about 1700 K),such as Co, Ni, and Fe, can exhibit a rapid grain growth over a moderatetemperature range of about 220-450° C. (about 493-723 K), resulting inmicrometer-sized grains at a temperatures less than a half of theirmelting temperatures. Approaches that involve adding solute atoms tostabilize nanocrystalline metals by pinning the grain boundaries orlowering the grain boundary energy have yielded some successful results.However, these approaches remain inherently constrained by thethermodynamic properties of the particular alloy systems. It is proposedthat dispersed nanoparticles could stabilize the UFG, nanocrystallinemetals at elevated temperatures. To evaluate the nanoparticle assistedstabilization in the as-cast Cu-13 WC nanocomposites at elevatedtemperatures, the grain growth was further investigated by STEM equippedwith in-situ heating capability. The in-situ heating path is staying atabout 200, about 400, and about 600° C. for about 10 min and about 850°C. for about 30 min. FIG. 4A and 4E show the local microstructures of assolidified Cu-13 WC samples with different local volume percentages ofnanoparticles. The local nanoparticle concentration in FIG. 4A is higherthan that in FIG. 4E. The relatively bright phases are WC nanoparticlesand the black phase is the Cu matrix. The grain structures at about 400,about 600 and about 850° C. shown in FIG. 4(B-D) and FIG. 4(F-H) areacquired at longer camera lengths than FIG. 4A and 4E to enhance thediffraction contrast of the image in order to distinguish individual Cugrains. In the nanoparticle rich local area, FIG. 4(A-D) show thatnanoscale grains were thermally stable at temperatures up to about 850°C. (FIG. 4D). In the area with fewer nanoparticles (FIG. 4(E-H)),ultrafine/nanoscale grains were thermally stable up to about 600° C.However, when the temperature was raised to about 850° C.,ultrafine/nanoscale grains started to grow and became larger grains(FIG. 4H). Although it is observed that the grains started to grow atabout 850° C., this growth was restricted by nanoparticles close to thisarea and therefore the grain sizes did not exceed the ultrafine grainsize. This restriction indicates that nanoparticles effectively inhibitgrain growth in solid state up to a temperature of about 850° C. (about1123 K), which significantly corresponds to about 0.83 of the meltingtemperature of Cu (about 1353 K).

To further validate the thermal stability of the refined Cu grains andevaluate the effects on mechanical properties by heat treatment, heatingis performed on the Cu-34 vol. % WC sample obtained by casting under aircooling (cooling rate of about 7 K/s), which had an average grain sizeof 208 ±94 nm, to about 750° C. (about 1023 K, about 0.75 of the meltingpoint of Cu) and the sample is held at this temperature for about 2.0hours under the protection of argon. FIG. 4I shows a typical SEM imageof the Cu-34 vol. % WC sample after the high temperature exposure. EBSDwas utilized to evaluate the potential grain growth in thenanoparticle-poor zone such as the marked white rectangle area in FIG.41. FIG. 4J clearly shows that most Cu grains are still smaller thanabout 1.0 μm while a significant number of nanoscale grains remainedafter the heat treatment. The grain size distribution in FIG. 4Kconfirms that the average Cu grain size is 248±135 nm after the heattreatment, which is close to the average grain size, 208±94 nm, beforethe heat treatment.

The thermal stability may be attributable to a Zener pinning effectderived from the presence of dispersed nanoparticles. A maximum meangrain size can be theoretically estimated by:

$D_{\max} = \frac{4r}{3f_{v}}$

where r is the nanoparticle radius and f_(v) is the volume fraction ofthe nanoparticles. The D_(max) for Cu-34% WC is estimated to be about392 nm, which is comparable with the experimental data. When comparedwith other precipitate particles from alloy systems, the ex-situ WCnanoparticles will not dissolve in the Cu matrix at high temperatures,offering a superior thermal stability than alloys.

Mechanical Properties

The presence of fine, well-dispersed nanoparticles in a metal matrix cansignificantly enhance mechanical responses. To gain fundamental insightsinto the property enhancement induced by the refined grains andnanoparticles, micropillar compression tests are conducted as shown inFIG. 11. The typical micropillar of Cu-13 WC sample with diameter ofabout 4 μm and length of about 9 μm is shown FIG. 11A, demonstratingthat dense nanoparticles are dispersed in the micropillar. Thepolycrystalline nature of the micropillar is shown in FIG. 11B. Thetypical results from the micropillar compression tests are shown in FIG.11C. The yield point increased from 180±8 MPa to 827±74 forsubstantially pure Cu and as-solidified Cu-34 vol. % WC, respectively.Moreover, micropillars from Cu-34 vol. % WC sustained a graduallyincreasing load up to about 1490 MPa and strain over about 25% withoutcatastrophic failure. In comparison, the substantially pure Cu sampleexperienced extensive slip as the strain increases discontinuously.Furthermore, it is observed that the mechanical properties did notchange much (Yield point drops from 827±74 to 780±9) after a thermalexposure at about 750° C. for about 2.0 hours, which further confirmsthat the Cu samples with WC nanoparticles exhibit excellent thermalstability. The detailed strengthening mechanisms are analyzed in thesupplementary subsection ‘Strengthening mechanisms’. The Young's modulusof substantially pure Cu, Cu-19 vol. % WC, and Cu-34 vol. % WC aremeasured by the nanoindentation method with a Berkovich tip. The Young'smodulus of pure Cu, Cu-19 vol. % WC, and Cu-34 vol. % WC are 139±8 GPa,192±13 GPa, and 223±16 GPa, respectively. It is hypothesized that theenhancement of Young's modulus is due to the high Young's modulus of WC(about 530-700 GPa) and the effective load transfer by thenanoparticles.

CONCLUSIONS

In summary, an approach to effectively control nucleation and graingrowth down to ultrafine/nanoscale during slow solidification isproposed by use of dispersed nanostructures in molten metals. Thenanostructures assisted grain refinement mechanisms, which combinecontinuous nucleation and grain growth control, and break thefundamental constraint of other grain refinement methods. This approachprovides a pathway to directly cast bulk UFG, nanocrystalline metalsunder slow cooling rates (e.g., less than about 100 K/s) for large scaleproduction. An unprecedented thermal stability up to about 1023 K, about0.75 T_(m) of Cu, for the as-cast bulk UFG, nanocrystalline Cu isdemonstrated. As-cast bulk UFG, nanocrystalline metals withnanoparticles also show exceptional strengths and Young's modulusenhancement. Furthermore, this general approach is applicable fordifferent materials systems such as Al-TiB₂ and Zn—WC. This approach canhave a significant impact on the solidification processes for metals,and also on numerous applications such as biomedical, chemical, andatmospheric sciences.

Materials and Methods: Materials Fabrication 1. Salt AssistedIncorporation Method

WC nanoparticles were mixed with Borax (Na₂B₄O₇)-5 wt.% CaF₂ saltpowders by a mechanical shaker (SK-0330-Pro) for about 1.0 h. The volumefraction of nanoparticles in the salt mixture is designed as about 10%.As shown in FIG. 5A, substantially pure oxygen free Cu (about 99.99%,Rotometals, Inc.) ingots were melted at about 1250° C. in a graphitecrucible by induction heater. Inert argon (Ar) gas was purged on themolten Cu to avoid severe oxidation. The mixed Na₂B₄O₇-5 wt.% CaF₂-WCnanoparticles were manually loaded on the surface of the molten Cu. Agraphite propeller was located below the Cu-salt interface and stirredat a speed of about 400 rpm for about 20 min to incorporate WCnanoparticles into Cu melt. Then the melt was allowed to cool down toabout 900° C. to allow Cu solidify first while the salt mixture wasstill in a liquid state. Liquid salt was poured out from the crucibleand leave a Cu-13 WC ingot. The volume fraction of WC nanoparticles inCu was designed to be 0, about 5, about 10, and about 20 vol. %.

Surface clean TiB₂ nanoparticles were synthesized by the magnesiothermicreduction of TiO₂ nanoparticles and B₂O₃ powders in molten salt. Thesynthesized TiB₂ nanoparticles and KAl₄ flux were then mechanicallymixed in solid state for about 3.0 h. Mixed powders were dehydrated atabout 120° C. for about 1.0 hour in a vacuum oven. An electricalresistance furnace was used to melt the Al ingots at about 820° C. underAr gas protection. Then, the mixed powders were added to the meltsurface and the melt was mechanically stirred at about 200 rpm for about10 min with titanium (Ti) mixing blade. The designed volume fraction ofTiB₂ in Al is about 10 vol. %. The melt was taken out from the furnaceand naturally cooled down to room temperature under Ar gas protection.

2. Powder Melting Method

Cu powders (Sigma Aldrich,<about 10 μm) and Zn powders (Alea Aesa, about150 μm) were first mixed with a designed volume fraction of WCnanoparticles (US Research Nanomaterials, about 150-200 nm) for about1.0 hour by a mechanical shaker (SK-0330-Pro). The mixed powders werethen compressed into disks with about 2.0 cm in diameter under about 250MPa by a hydraulic machine. As shown in FIG. 5B, Cu-13 WC cold compacteddisks were compression-melted at about 1250° C. by an induction heaterwhile under a pressure of about 7-10 MPa under a graphite disk andalumina piston. After solidification in air, Cu ingots with WCnanoparticles were obtained. Zn—WC cold compacted billets were thenmelted at about 500° C. and ultrasonic processed by a Niobium (Nb) probefor about 10 min to disperse WC nanoparticles in Zn melt. Then the meltwas taken out from the furnace and cooled down under Ar gas protection.

Cooling Rates Measurement

To evaluate the effects of the cooling rate, Cu-13 WC samples were thenmelted again at about 1250° C. under a protection of Argon gas and thencooled down in furnace, air and water. The cooling rate was measured bya K-type thermocouple connected to an Arduino UNO board for recordingthe cooling curve. The measured cooling rates of furnace cooling, aircooling and water quenching are about 2-4 K/s, about 7-12 K/s and about70-100 K/s, respectively. The typical cooling curves are shown in FIG.6, and the thermal arrest caused by the solidification of Cu is clearlyidentified in the furnace and air cooling curve, while not so noticeablein the water quenching curve due to the rapid heat dissipation.

Structure Characterization

The microstructure, distribution and dispersion of nanoparticles inmetals were evaluated by scanning electron microscope (SEM), focus ionbeam (FIB) imaging, electron backscatter diffraction (EBSD) andtransmission electron microscope (TEM). To clean the surface and revealthe nanoparticles in metal matrix, the mechanically ground and polishedas-cast samples were further polished by low-angle ion milling (ModelPIPS 691, Gatan). SEM images were acquired at 0° and a tilted about 52°by ZEISS Supra 40VP and FEI Nova 600, respectively. The composition ofthe material was characterized by energy-dispersive X-ray spectroscopy(EDS).

The volume fraction of nanoparticles was estimated based on the atomicfraction of the major element in the base metal and nanoparticles.Taking advantages of the channeling contrast of different grains inducedfrom the ion beam, FIB imaging was used to reveal the grain structure.Grain size and orientation were evaluated by EBSD (FEI Quanta 3D) atabout 30 kV with a current of about 12 nA and FIB with a channeldetection electron multiplier (CDEM) detector. The interfaces betweenmatrix and nanoparticles were evaluated by a FEI Titan TEM at about 300kV. The thin film TEM samples were machined by FIB.

Solidification Behavior

To evaluate the solidification behaviors during cooling, Cu-13 WC andsubstantially pure Cu samples with the same mass (about 50 mg) wereanalyzed by Differential Scanning calorimetry (DSC) (TA Instruments,Q600). Samples were heated up to about 1250° C. with a heating rate ofabout 40° C./min and then cooled down with a cooling rate of about 5°C./min.

High Temperature Stability

To evaluate the UFG, nanocrystalline structures stability at hightemperatures, in situ heating Scanning Transmission Electron Microscope(STEM) was conducted at about 300 kV with a convergence angle of about4.3 mrad and geometric aberrations in the probe corrected to thirdorder. The small convergence angle was chosen to enhance diffractioncontrast for grain identification. An FIB milled TEM sample was quicklyheated to about 200° C., about 400° C., about 600° C., about 700° C. andabout 850° C. by a Gatan in-situ heating holder and held for about 10min at each step for imaging. The diffraction camera length wasincreased at higher temperatures in order to enhance diffractioncontrast, emphasizing the difference between individual Cu grains forrobust identification. The collection semi-angle was about 48-240 mradat room temperature, about 15-75 mrad at about 400° C., and about 12-60mrad at about 600° C. and above. Moreover, the bulk solidified Cu-34vol. % WC (air cooling, about 7 K/s) samples were heated up to about750° C. (about 1023K, about 0.75 of Tm) and stayed at that temperaturefor about 2.0 hours under the protection of Ar gas. The grain structurewas then evaluated by FIB imaging and EBSD. Mechanical properties wereevaluated by microcompression test.

Mechanical characterization and properties

An MTS nanoindenter with a flat punch tip was used for microcompressiontests at a strain rate of about 5×10⁻² s⁻¹ under room temperature.Micropillars of about 4 μm in diameter and about 9 μm in length weremachined by FIB. To evaluate the elastic modulus, microindentation testswith an indent depth of about 2 μm were performed by the same MTSnanoindenter with a Berkovich tip. For each sample, at least ten pointswere measured. Accordingly micropillars with a diameter and a length ofabout 4 μm and about 9 μm, respectively, were machined by FIB from theas-solidified (air cooling, cooling rate of about 7 K/s) and heattreated samples (about 750° C. for about 2.0 hours).

Nanoparticle Dispersion and Self-Stabilization Mechanism

Silicon carbide nanoparticles can be dispersed in magnesium matrix by aself-dispersion and self-stabilization mechanism. There are three majorfactors contributing in this mechanism: (1) Good wetting between moltenmetal and nanoparticles creates an energy barrier to prevent atomiccontact and sintering of nanoparticles in the melt. In the Cu-13 WCsystem, the energy barrier can be calculated by the following equation:

W_(barrier)=Sσ_(Cu) cos θ

where S is the effective area and can be calculated by S=πRD₀ (D₀=0.2nm), _(σCu) is the surface energy of Cu at the processing temperature(about 1.27 J m⁻²), and θ is the wetting angle, e.g. at about 1250° C.(about)10° . In the Cu-13 WC system, an energy barrier of 7×10⁴ zJ isobtained; (2) Thermal energy that allow nanoparticles to move randomlyin the molten melt by overcoming the attractive van der Waals potentialbetween nanoparticles. A higher thermal energy is desired. Theprocessing temperature for Cu-13 WC is about 1250° C., which provide athermal energy of E=k_(b)T=21.0 zJ. (3) A van der Waals potentialbetween nanoparticles to lessen the attraction of nanoparticles fromeach other to form nanoparticles clusters in molten metals. A smallattractive van der Waals potential is desired. It can be calculated bythe following equation:

$W_{vdw} = {{- \frac{\left( {\sqrt{A_{Cu}} - \sqrt{A_{WC}}} \right)^{2}}{6D}}\frac{R}{2}}$

where A_(Cu=)410 zJ and A_(WC) are the Hamaker constants, D is thedistance between two nanoparticles that can be as small as two atomiclayer thick (about 0.4 nm), R is the radius of the nanoparticle (about100 nm). Although the data of A_(wc) is not available, a 7×10⁴ zJ energybarrier is several orders of magnitude higher than thermal energy. SinceWC is a conductive ceramic material, A_(wc) is estimated to be in therange of 200 to 500 zJ, indicating that W_(barrier) would always be muchhigher than W_(vdw) for stabilization of dispersed WC nanoparticles inCu melt.

Growth Restriction Factor by Nanoparticles

For growth restriction by extra solute atoms, the constitutionalundercooling can be described by:

${\Delta Tc} = {m{C_{0}\left( {1 - \frac{1}{\left( {1 - f_{s}} \right)^{({1 - k})}}} \right)}}$

where ΔT_(c) is the constitutional undercooling, m is the liquidus slopein a linear phase diagram, C₀ is the solute content in the alloy and kis the equilibrium solute partition coefficient, and f_(s) is the solidfraction solidified. By taking the derivative of this equation respectto f_(s), Q can be expressed as:

$Q = {{m{C_{0}\left( {k - 1} \right)}} = \left( \frac{\partial\left( {\Delta T_{c}} \right)}{\partial f_{s}} \right)_{{fs}\rightarrow 0}}$

This equation indicates that the physical meaning of growth restrictionfactor is the initial rate of development of constitutionalundercooling. Taking Al—Ti as one example, the constitutionalundercooling of Al—Ti alloy system is shown in FIG. 12. Ti is the one ofthe most effective atoms to restrict grain growth in Al. The growthrestriction factor is described as Q₁, Q₂ and Q₃ for different Ticoncentration of 0.05, 0.10 and 0.15, respectively. With higherconcentration of Ti atoms in Al melt, the growth restriction factorincreases. As marked by the black arrow in the FIG. 12, a steeper slopeprovides larger growth restriction factor. However, the maximumsolubility of Ti atoms in Al is 0.15 which constrains Q to reach amaximum value as Q_(max).

In contrast, the growth restriction factor, Q_(np), introduced bynanoparticle assisted phase control break the fundamental constraint setby Q that depends on constitutional undercooling. As far as thesolidification front touches with a nanoparticle with curved shape or ananoscale channel between nanoparticles, unlike the constitutionalundercooling built gradually ahead of the solidification front, anundercooling is immediately established by the Gibbs-Thompson effect. Asshown in FIG. 12, the undercooling profile established by Gibbs-Thompsoneffect is a step function with an infinitely large slope at the initialpoint. Thus Q_(np) which is the initial rate of the undercoolingdevelopment by nanoparticles, can be readily increased to asignificantly large number, if not infinity.

Strengthening Mechanism

The strengthening from nanoparticles and refined grain structures in theas-solidified samples of Cu-34 vol. % WC is about 647 MPa. The majorstrengthening mechanisms include Orowan strengthening induced from thepopulous and dispersed nanoparticles, Hall-Petch effect, and loadbearing transfer.

The contribution from Orowan strengthening (Δσ_(Orowan) ) fromwell-dispersed nanoparticles can be calculated by the followingequation:

${\Delta \sigma_{Orowan}} = {\frac{{0.1}3G_{m}b}{d_{p}\left\lbrack {\left( \frac{1}{2V_{p}} \right)^{\frac{1}{3}} - 1} \right\rbrack}\ln \frac{d_{p}}{2b}}$

where G_(m), b, V_(p) and d_(p) are the shear modulus of the matrix, theBurger vector, the volume fraction and the size of the nanoparticles,respectively. In this example, G_(m)=46 GPa, b=0.256 nm, V_(p)=0.34 andd_(p)=200 nm, and the calculated Δσ_(Orowan) is 333 MPa. It should benoted that this value is estimated based on ideal dispersion.

The contribution from Hall-Petch effect can be calculated by thefollowing equation:

Δσ_(y)=kd^(−1/2)

where d is the grain size and k is a material constant. For Cu, k=0.11MPa·m, and d=208 nm, and the calculated Δσ_(y)=246 MPa.

The load bearing strengthening can be calculated by the followingequation:

Δσ_(load)=1.5V_(p)σ_(i)

where σ_(i) is the interfacial bonding strength between nanoparticlesand metal matrix. The strong interfacial bonding between Cu and WCnanoparticles contributes in the load bearing mechanism. It can bedifficult to estimate the interfacial bonding strength.

Definitions

As used herein, the singular terms “a,” “an,” and “the” may includeplural referents unless the context clearly dictates otherwise. Thus,for example, reference to an object may include multiple objects unlessthe context clearly dictates otherwise.

As used herein, the term “set” refers to a collection of one or moreobjects. Thus, for example, a set of objects can include a single objector multiple objects.

As used herein, the terms “connect,” “connected,” and “connection” referto an operational coupling or linking. Connected objects can be directlycoupled to one another or can be indirectly coupled to one another, suchas via one or more other objects.

As used herein, the terms “substantially” and “about” are used todescribe and account for small variations. When used in conjunction withan event or circumstance, the terms can refer to instances in which theevent or circumstance occurs precisely as well as instances in which theevent or circumstance occurs to a close approximation. When used inconjunction with a numerical value, the terms can refer to a range ofvariation of less than or equal to ±10% of that numerical value, such asless than or equal to ±5%, less than or equal to ±4%, less than or equalto ±3%, less than or equal to ±2%, less than or equal to ±1%, less thanor equal to ±0.5%, less than or equal to ±0.1%, or less than or equal to±0.05%. For example, a first numerical value can be “substantially” or“about” the same as or equal to a second numerical value if the firstnumerical value is within a range of variation of less than or equal to±10% of the second numerical value, such as less than or equal to ±5%,less than or equal to ±4%, less than or equal to ±3%, less than or equalto ±2%, less than or equal to ±1%, less than or equal to ±0.5%, lessthan or equal to ±0.1%, or less than or equal to ±0.05%.

As used herein, the term “size” refers to a characteristic dimension ofan object. Thus, for example, a size of an object that is spherical orspheroidal can refer to a diameter of the object. In the case of anobject that is non-spherical or non-spheroidal, a size of the object canrefer to a diameter of a corresponding spherical or spheroidal object,where the corresponding spherical or spheroidal object exhibits or has aparticular set of derivable or measurable properties that aresubstantially the same as those of the non-spherical or non-spheroidalobject. When referring to a set of objects as having a particular size,it is contemplated that the objects can have a distribution of sizesaround the particular size. Thus, as used herein, a size of a set ofobjects can refer to a typical size of a distribution of sizes, such asan average size, a median size, or a peak size.

As used herein, the term “nanostructure” refers to an object that has atleast one dimension in a range of about 1 nm to about 1000 nm. Ananostructure can have any of a wide variety of shapes, and can beformed of a wide variety of materials. Examples of nanostructuresinclude nanofibers, nanoplatelets, and nanoparticles.

As used herein, the term “nanoparticle” refers to a nanostructure thatis generally or substantially spherical or spheroidal. Typically, eachdimension of a nanoparticle is in a range of about 1 nm to about 1000nm, and the nanoparticle has an aspect ratio of about 5 or less, such asabout 3 or less, about 2 or less, or about 1.

As used herein, the term “nanofiber” refers to an elongatednanostructure. Typically, a nanofiber has a lateral dimension (e.g., awidth) in a range of about 1 nm to about 1000 nm, a longitudinaldimension (e.g., a length) in a range of about 1 nm to about 1000 nm orgreater than about 1000 nm, and an aspect ratio that is greater thanabout 5, such as about 10 or greater.

As used herein, the term “nanoplatelet” refers to a planar-like,nanostructure.

Additionally, amounts, ratios, and other numerical values are sometimespresented herein in a range format. It is to be understood that suchrange format is used for convenience and brevity and should beunderstood flexibly to include numerical values explicitly specified aslimits of a range, but also to include all individual numerical valuesor sub-ranges encompassed within that range as if each numerical valueand sub-range is explicitly specified. For example, a ratio in the rangeof about 1 to about 200 should be understood to include the explicitlyrecited limits of about 1 and about 200, but also to include individualratios such as about 2, about 3, and about 4, and sub-ranges such asabout 10 to about 50, about 20 to about 100, and so forth.

While the disclosure has been described with reference to the specificembodiments thereof, it should be understood by those skilled in the artthat various changes may be made and equivalents may be substitutedwithout departing from the true spirit and scope of the disclosure asdefined by the appended claim(s). In addition, many modifications may bemade to adapt a particular situation, material, composition of matter,method, operation or operations, to the objective, spirit and scope ofthe disclosure. All such modifications are intended to be within thescope of the claim(s) appended hereto. In particular, while certainmethods may have been described with reference to particular operationsperformed in a particular order, it will be understood that theseoperations may be combined, sub-divided, or re-ordered to form anequivalent method without departing from the teachings of thedisclosure. Accordingly, unless specifically indicated herein, the orderand grouping of the operations are not a limitation of the disclosure.

What is claimed is:
 1. A nanocrystalline material comprising: a matrixincluding one or more metals; and nanostructures dispersed in thematrix, wherein the matrix is polycrystalline and includes grains havingan average size of about 1 μm or less.
 2. The nanocrystalline materialof claim 1, wherein the average size of the grains is about 600 nm orless.
 3. The nanocrystalline material of any of claim 1, wherein theaverage size of the grains is about 400 nm or less.
 4. Thenanocrystalline material of any of claim 1, wherein the nanostructuresare dispersed in the matrix at a volume fraction of about 5% or greaterof the nanocrystalline material.
 5. The nanocrystalline material ofclaim 4, wherein the volume fraction of the nanostructures in thenanocrystalline material is about 10% or greater.
 6. The nanocrystallinematerial of claim 4, wherein the volume fraction of the nanostructuresin the nanocrystalline material is about 15% or greater.
 7. Thenanocrystalline material of claim 1, wherein the matrix includes copper,and the nanostructures include a transition metal or a transition metalcarbide.
 8. The nanocrystalline material of claim 1, wherein the matrixincludes zinc, and the nanostructures include a transition metal or atransition metal carbide.
 9. The nanocrystalline material of claim 1,wherein the matrix includes aluminum, and the nanostructures include atransition metal carbide or a transition metal boride.
 10. Amanufacturing method of a nanocrystalline material, comprising: heatinga matrix material including one or more metals to form a melt; loading amixture including a salt and nanostructures over a surface of the melt,such that the salt is heated to form a molten salt including thenanostructures dispersed therein; agitating the melt to incorporate thenanostructures from the molten salt into the melt; and cooling the meltincluding the nanostructures dispersed therein to form a mastermaterial.
 11. The manufacturing method of claim 10, further comprisingsubjecting the master material to casting to form a metal part.
 12. Themanufacturing method of claim 11, wherein subjecting the master materialto casting includes heating the master material to form a mastermaterial melt, delivering the master material melt to a mold defining ahollow space with a requisite shape, and cooling and solidifying themaster material melt to form the metal part including thenanocrystalline material and having the requisite shape.
 13. Themanufacturing method of claim 12, wherein cooling the master materialmelt is performed at a rate of less than about 100 K/s.
 14. Amanufacturing method of a nanocrystalline material, comprising: mixing apowder of a matrix material and reinforcing nanostructures to form apowder mixture; compacting the powder mixture to form a preform; heatingthe preform under compression to form a melt including thenanostructures dispersed therein; and cooling the melt including thenanostructures dispersed therein to form a master material.
 15. Themanufacturing method of claim 14, further comprising subjecting themaster material to casting to form a metal part.
 16. The manufacturingmethod of claim 15, wherein subjecting the master material to castingincludes heating the master material to form a master material melt,delivering the master material melt to a mold defining a hollow spacewith a requisite shape, and cooling and solidifying the master materialmelt to form the metal part including the nanocrystalline material andhaving the requisite shape.
 17. The manufacturing method of claim 16,wherein cooling the master material melt is performed at a rate of lessthan about 100 K/s.